Development of Ni-Sr(V,Ti)O3-δ Fuel Electrodes for Solid Oxide Fuel Cells

A series of strontium titanates-vanadates (STVN) with nominal cation composition Sr1-xTi1-y-zVyNizO3-δ (x = 0–0.04, y = 0.20–0.40 and z = 0.02–0.12) were prepared by a solid-state reaction route in 10% H2–N2 atmosphere and characterized under reducing conditions as potential fuel electrode materials for solid oxide fuel cells. Detailed phase evolution studies using XRD and SEM/EDS demonstrated that firing at temperatures as high as 1200 °C is required to eliminate undesirable secondary phases. Under such conditions, nickel tends to segregate as a metallic phase and is unlikely to incorporate into the perovskite lattice. Ceramic samples sintered at 1500 °C exhibited temperature-activated electrical conductivity that showed a weak p(O2) dependence and increased with vanadium content, reaching a maximum of ~17 S/cm at 1000 °C. STVN ceramics showed moderate thermal expansion coefficients (12.5–14.3 ppm/K at 25–1100 °C) compatible with that of yttria-stabilized zirconia (8YSZ). Porous STVN electrodes on 8YSZ solid electrolytes were fabricated at 1100 °C and studied using electrochemical impedance spectroscopy at 700–900 °C in an atmosphere of diluted humidified H2 under zero DC conditions. As-prepared STVN electrodes demonstrated comparatively poor electrochemical performance, which was attributed to insufficient intrinsic electrocatalytic activity and agglomeration of metallic nickel during the high-temperature synthetic procedure. Incorporation of an oxygen-ion-conducting Ce0.9Gd0.1O2-δ phase (20–30 wt.%) and nano-sized Ni as electrocatalyst (≥1 wt.%) into the porous electrode structure via infiltration resulted in a substantial improvement in electrochemical activity and reduction of electrode polarization resistance by 6–8 times at 900 °C and ≥ one order of magnitude at 800 °C.


Introduction
Concerns about the disadvantages of traditional Ni-YSZ (YSZ = yttria-stabilized zirconia) anodes of solid oxide fuel cells (SOFCs) have boosted significant research activities on the development of alternative anode materials [1,2]. Ni-YSZ cermets have poor redox stability, and significant volume changes upon occasional nickel oxidation/re-reduction cycles cause irreversible microstructural degradation [3][4][5]. Nickel agglomeration is another common degradation mechanism that leads to a loss of electrical connectivity between Ni particles and a decrease in electrocatalytic activity [5,6]. In addition, Ni-YSZ anodes are readily poisoned by sulfur and suffer from carbon deposition, which are key issues for direct operation with hydrocarbon fuels [7,8].
Nominal cation composition, perovskite lattice parameters, and density of Sr 1-α Ti 1-β(1+γ) V β Ni βγ O 3-δ (STVN) ceramics. Mechanically activated precursor powders were subjected to thermal treatments in a controlled atmosphere of 10 vol.% H 2 in N 2 . In order to assess the effects of temperature and time of thermal treatment on the phase formation, powdered samples were calcined at different temperatures between 800 and 1200 • C for 4-10 h, with intermediate regrinding in a mortar with ethanol when appropriate. The final powders were prepared by calcination of precursor mixtures at 1200 • C for 30 h in total (see discussion below). These as-synthesized powders were ball-milled with ethanol for 4 h at 150 rpm, dried, and then used for the preparation of ceramic samples and electrodes. The disk-shaped ceramic samples were compacted by uniaxial pressing (Ø 18 mm, thickness~2 mm, 40 MPa) followed by isostatic pressing (230 MPa) and sintered in flowing 10% H 2 -N 2 at 1450 • C for 10 h.
Sintered ceramic pellets were polished using SiC grinding paper (Buehler, Leinfelden-Echterdingen, Germany). The density of the prepared ceramics (ρ exp ) was calculated from the geometric dimensions and masses of the polished samples. Rectangular bars (approximate dimensions: 1.5 × 2.5 × 12 mm) for electrical and dilatometric measurements were cut out of disk-shaped samples using a Struers Minitom precision cutting machine (Struers, Copenhagen, Denmark) with a diamond cut-off wheel. Powdered samples for X-ray diffraction (XRD) were prepared by grinding sintered ceramics in a mortar.

Characterization of Materials
XRD patterns of powdered samples were recorded using a PANalytical X'pert PRO diffractometer (CuK α radiation, 2Θ range of 10-80 • , PANalytical, Almelo, The Netherlands). Phase identification and quantification were carried out using X'Pert Highscore Plus software (PANalytical). The unit cell parameters were calculated using FullProf software (v.7.20, ILL, Grenoble, France) employing the profile matching method. Microstructural characterization of STVN ceramics was performed through scanning electron microscopy (SEM) using a Hitachi SU-70 microscope (Hitachi, Tokyo, Japan) equipped with a Bruker Quantax 400 detector (Bruker, Berlin, Germany) for energy dispersive spectroscopy (EDS) analysis. Thermal expansion of STVN ceramics between 25 and 1100 • C was studied using controlled-atmosphere dilatometry in a vertical Linseis L75 dilatometer on heating/cooling (3 • C/min) in a flowing 10% H 2 -N 2 mixture.
The electrical conductivity (σ) of the STVN ceramics was measured as a function of temperature and oxygen partial pressure p(O 2 ) using the four-probe DC technique with bar-shaped samples. Pt wires were used as probes and current collectors. Additionally, the end-face surfaces of the bars were covered with Pt paint (Heraeus CL11-5349, Hanau, Germany) to improve the electrical contact. DC was supplied by a Yokogawa 7651 source and the voltage between the potential probes was measured using an Agilent 34460A multimeter. The experiments were performed in a cooling regime in a flowing 10% H 2 -N 2 atmosphere with stepwise decreases in temperature from 1000 to 300 • C and isothermally with a stepwise increase in p(O 2 ) from 10 −20 to 10 −16 atm. The p(O 2 ) was controlled by the composition of 10% (H 2 + H 2 O)-N 2 gas mixtures and measured using an electrochemical 8 mol.% yttria-stabilized zirconia oxygen sensor. The minimum equilibration time before each measurement was 30 min.

Fabrication and Electrochemical Characterization of Electrodes
The electrochemical activity of the STVN electrodes was evaluated using an electrolytesupported symmetrical cell configuration. For the fabrication of solid electrolyte membranes, commercial yttria-stabilized zirconia (ZrO 2 ) 0.92 (Y 2 O 3 ) 0.08 (8YSZ, Tosoh, Tokyo, Japan) powders were compacted uniaxially into disks (Ø 15 mm), sintered in air at 1600 • C for 10 h, and then polished down to a thickness of 1.3 mm. STVN inks were prepared using as-synthesized powders (30 vol.%) mixed with an organic vehicle composed of α-terpineol (solvent) and ethyl cellulose (binder) with additions of stearic acid as a dispersant agent (extra 3 wt.%). The ink components were blended in a planetary ball mill (Retsch S1 mill, nylon vial, tetragonal zirconia balls) for 4 h at 150 rpm. Circular electrodes with an effective area of 0.2 cm 2 were brush-painted symmetrically on each side of the 8YSZ solid electrolyte disk and sintered in a flowing 10% H 2 -N 2 gas mixture at 1100 • C for 2 h (heating/cooling rate of 3 • C/min) to consolidate the porous electrode layers with~40-50 µm thickness.
The performance of the symmetrical half-cells was evaluated by electrochemical impedance spectroscopy (EIS) using an AUTOLAB PGSTAT 302 instrument (Eco Chemie, Utrecht, Netherlands) equipped with a frequency response analyzer (FRA2) module. Pt gauzes with Pt wires were used as current collectors. The experiments were performed in a humidified fuel atmosphere (~3% H 2 O-10% H 2 -N 2 ) under zero-DC conditions. The EIS spectra were registered in the frequency range from 0.01 Hz to 1 MHz with an AC amplitude of 50 mV. After heating, the cell was held at 900 • C for ≥20 h before starting the experiment. Following that, the measurements were carried out in a cooling regime from 900 to 700 • C with a 50 • C step. The stabilization time at each step was~1 h. A post-mortem SEM/EDS inspection was performed in order to examine the microstructure of the STVN electrodes after the EIS measurements.

Electrode Modifications by Infiltration
The modification of STVN electrodes with gadolinia-doped ceria Ce 0.9 Gd 0.1 O 2-δ (CGO) and metallic Ni was conducted via the infiltration procedure. Several aqueous solutions of metal nitrates were prepared using Ce(NO 3 ) 3 ·6H 2 O (Sigma-Aldrich, ≥99.9%), Gd(NO 3 ) 3 ·6H 2 O (Alfa Aesar, ≥99.9%), and Ni(NO 3 ) 2 ·6H 2 O (Sigma-Aldrich, ≥99.9%). The solutions contained metal cations in different proportions: (i) Ce:Gd = 9:1, (ii) Ce:Gd:Ni = 9:1:1, and (iii) Ce:Gd:Ni = 9:1:5. Careful infiltration of a solution into the porous electrode structure was followed by drying and thermal treatment at 900 • C in a 10% H 2 -N 2 atmosphere for 2 h for decomposition of nitrates and formation of dispersed CGO and Ni particles. The load of infiltrated components was calculated from the change in the weight of the cells. When necessary, the infiltration process was repeated to achieve the target load in each electrode (~20-30 wt.%).

Effect of Mechanochemical Treatment
The mechanical activation of the precursor mixtures resulted in a decrease in the intensity and a noticeable broadening of the main XRD reflections ( Figure S1). This could be caused by a number of factors, including the reduction in the particle size, the lattice strain induced by high-energy milling, and even partial amorphization of initial crystalline phases. The onset of the target SrTiO 3 -based perovskite phase was observed for some compositions ( Figure S1). Overall, the mechanochemical treatment resulted in a rather limited output in terms of the formation of the desired product but provided very good precursor homogenization, which was favorable for further synthetic process.

Effect of Calcination Temperature
The mechanically-activated precursor mixtures were calcined for 4 h in a flowing 10% H 2 -N 2 mixture at three different temperatures (800, 900, and 1000 • C) with the aim to find optimal conditions for the easy formation of the target perovskite phase. Figure 1 shows the XRD pattern of calcined samples of two selected compositions: S100V40N12 without A-site vacancies and with the highest nominal nickel content and S96V30N3 with A-site vacancies and low nickel content. The thermal treatment at 800 • C was not very effective; the corresponding XRD patterns revealed the presence of substantial amounts of secondary products, including Sr 3 (VO 4 ) 2 , metallic Ni, traces of unreacted TiO 2 , and other phases. Increasing the temperature of calcination up to 900 • C substantially promoted the formation of the SrTiO 3 -based perovskite that became the main phase after this treatment. Nonetheless, all compositions still comprised substantial amounts of secondary phases including Sr 3 (VO 4 ) 2 (particularly in A-site stoichiometric compositions), the second perovskite phase based on SrVO 3 , and metallic Ni. The calcination at 1000 • C yielded very similar results, with only a minor decrease in the fraction of Sr 3 (VO 4 ) 2 impurity. The XRD patterns and estimated fractions of different phases for all STVN compositions after calcinations at 800-1000 • C can be found in Figures S2-S4. Analysis of the results revealed the correlations between the nominal composition and the fractions of secondary phases. The amount of segregated metallic nickel phase is rationally interrelated with the nominal nickel content. The fraction of Sr 3 (VO 4 ) 2 impurity increases with the increasing nominal Ni content but decreases with the increasing nominal A-site deficiency. Increasing vanadium content promotes the segregation of SrVO 3 -based perovskite, particularly for nominally Sr-deficient compositions. Note that, due to the very close lattice parameters of SrTiO 3 and SrVO 3 phases and, consequently, the overlapping of the XRD reflections, the detection and quantification of SrVO 3 secondary phase was not always possible. In summary, the results of phase analysis imply that nickel does not tend to incorporate into the B-sublattice of Sr(Ti,V)O 3 perovskites during the thermal treatment of mechanically-activated precursors in reducing atmospheres at 800-1000 • C but preferentially segregates in metallic form. This results in an excess of strontium, particularly in the compositions with a higher nominal nickel content, which, in turn, promotes the formation of a strontium orthovanadate Sr 3 (VO 4 ) 2 phase. 1000 °C but preferentially segregates in metallic form. This results in an excess of strontium, particularly in the compositions with a higher nominal nickel content, which, in turn, promotes the formation of a strontium orthovanadate Sr3(VO4)2 phase. Figure 1. XRD patterns of S100V40N12 and S96V30N3 samples after calcination for 4 h in a flowing 10% H2-N2 atmosphere at 800, 900, and 1000 °C.

Phase Evolution in Consecutive Calcination Steps
As direct calcinations were not efficient for the preparation of phase-pure perovskitetype solid solutions, the evolution of the phase compositions was assessed in consecutive thermal treatment steps of pelletized samples in flowing 10% H2-N2 at 800-1100 °C. After each calcination, the samples were ground in a mortar with ethanol, compacted again, and subjected to the next step of thermal treatment.
The estimated fractions of different phases for the two selected compositions (S100V40N12 and S96V30N3) are shown in Figure 2. Corresponding XRD patterns can be found in Figures S5 and S6. The results of thermal treatments at lower temperatures, 800-900 °C, were qualitatively identical to those obtained by the direct calcinations of precursors. In the case of nominally Sr-stoichiometric S100V40N12 with a large Ni content, subsequent firing steps at 1000 °C/4h and 1100 °C/4h did not result in a noticeable improvement. The sample still comprised a large fraction of Sr3(VO4)2 (~30-33 wt.%), as

Phase Evolution in Consecutive Calcination Steps
As direct calcinations were not efficient for the preparation of phase-pure perovskitetype solid solutions, the evolution of the phase compositions was assessed in consecutive thermal treatment steps of pelletized samples in flowing 10% H 2 -N 2 at 800-1100 • C. After each calcination, the samples were ground in a mortar with ethanol, compacted again, and subjected to the next step of thermal treatment.
The estimated fractions of different phases for the two selected compositions (S100V40N12 and S96V30N3) are shown in Figure 2. Corresponding XRD patterns can be found in Figures S5 and S6. The results of thermal treatments at lower temperatures, 800-900 • C, were qualitatively identical to those obtained by the direct calcinations of precursors. In the case of nominally Sr-stoichiometric S100V40N12 with a large Ni content, subsequent firing steps at 1000 • C/4 h and 1100 • C/4 h did not result in a noticeable improvement. The sample still comprised a large fraction of Sr 3 (VO 4 ) 2 (~30-33 wt.%), as well as perovskitetype SrVO 3 in addition to the main SrTiO 3 -based perovskite phase and metallic nickel. It was noticed that the amount of undesired secondary phases decreased with the increasing calcination time at 1100 • C; however, this occurred very slowly. On the contrary, only a trace amount of Sr 3 (VO 4 ) 2 impurity was detected in the S96V30N3 sample after the thermal treatment step at 1000 • C, and no evidence of this secondary phase was found after subsequent calcination at 1100 • C. Nonetheless, the second perovskite phase was still present even after calcinations at 1100 • C, as suggested by the asymmetry of the reflections of the main perovskite phase ( Figure S6); its fraction did not alter with the treatment time ( Figure 2). well as perovskite-type SrVO3 in addition to the main SrTiO3-based perovskite phase and metallic nickel. It was noticed that the amount of undesired secondary phases decreased with the increasing calcination time at 1100 °C; however, this occurred very slowly. On the contrary, only a trace amount of Sr3(VO4)2 impurity was detected in the S96V30N3 sample after the thermal treatment step at 1000 °C, and no evidence of this secondary phase was found after subsequent calcination at 1100 °C. Nonetheless, the second perovskite phase was still present even after calcinations at 1100 °C, as suggested by the asymmetry of the reflections of the main perovskite phase ( Figure S6); its fraction did not alter with the treatment time ( Figure 2). Thus, consecutive thermal treatment steps did not lead to any noticeable improvements in terms of the formation of the desired phase. This appears to be in agreement with the literature data, which shows that V 5+ -based strontium orthovanadate Sr3(VO4)2 is a stable phase and forms as an intermediate in the course of reduction of oxidized strontium pyrovanadate Sr V O to reduced SrV 4+ O3 perovskite [16]. Typically, heat treatments at temperatures above 1000 °C are required to eliminate this phase [12,16]. In the present work, the dissolution of the intermediate secondary phases was found to be very slow even at 1100 °C.

Final Synthetic Procedure
Since calcinations at 1100 °C for up to 20 h were still insufficient for the formation of a single Sr(Ti,V)O3 perovskite phase combined with metallic nickel, it was decided to conduct solid-state synthesis at a temperature of 1200 °C, which was expected to be low enough to avoid excessive grain growth. Fresh mechanically activated precursors of all compositions were pelletized and fired at 1200 °C for 10 h. XRD inspection of the samples after this thermal treatment revealed that some compositions with a higher nominal Ni content combined with a low or no A-site deficiency (namely, S100V30N6, S100V40N12, and S98V30N9) still contained a non-negligible fraction of Sr3(VO4)2 impurity ( Figure S7). These three compositions were excluded from further work. The presence of the second Thus, consecutive thermal treatment steps did not lead to any noticeable improvements in terms of the formation of the desired phase. This appears to be in agreement with the literature data, which shows that V 5+ -based strontium orthovanadate Sr 3 (VO 4 ) 2 is a stable phase and forms as an intermediate in the course of reduction of oxidized strontium pyrovanadate Sr 2 V 5+ 2 O 7 to reduced SrV 4+ O 3 perovskite [16]. Typically, heat treatments at temperatures above 1000 • C are required to eliminate this phase [12,16]. In the present work, the dissolution of the intermediate secondary phases was found to be very slow even at 1100 • C.

Final Synthetic Procedure
Since calcinations at 1100 • C for up to 20 h were still insufficient for the formation of a single Sr(Ti,V)O 3 perovskite phase combined with metallic nickel, it was decided to conduct solid-state synthesis at a temperature of 1200 • C, which was expected to be low enough to avoid excessive grain growth. Fresh mechanically activated precursors of all compositions were pelletized and fired at 1200 • C for 10 h. XRD inspection of the samples after this thermal treatment revealed that some compositions with a higher nominal Ni content combined with a low or no A-site deficiency (namely, S100V30N6, S100V40N12, and S98V30N9) still contained a non-negligible fraction of Sr 3 (VO 4 ) 2 impurity ( Figure S7). These three compositions were excluded from further work. The presence of the second SrVO 3 -based perovskite phase was still suspected at least for one of the remaining compositions (S96V30N3), based on the asymmetry of the XRD reflections. Therefore, six STVN compositions were reground and fired in pelletized form at 1200 • C for an additional 20 h to ensure the full dissolution of secondary oxide phases ( Figure S8). These as-synthesized materials were powdered and then used for the preparation of electrode layers and ceramic samples.

Crystal Structure and Microstructure
XRD analysis confirmed that both as-synthesized STVN powders (1200 • C, 30 h) and sintered STVN ceramics (1450 • C, 10 h) comprised a Sr 1-a Ti 1-b V b O 3-δ solid solution with a cubic perovskite-type lattice (space group Pm3m) isostructural to the parent SrTiO 3 and SrVO 3 , and a fraction of metallic nickel phase. In agreement with the data on the Sr(Ti,V)O 3-δ system [34,48], the unit cell parameter of the perovskite phase shows a tendency to decrease with increasing vanadium content (Table 1), which is consistent with the values of ionic radii for six-fold coordinated V 4+ and Ti 4+ (0.58 Å and 0.605 Å, respectively [49]). The only exception to this trend is S96V40N8, the sample with the highest nickel content.
If nickel does not incorporate into the B sublattice of the perovskite structure and segregates completely as the metallic phase, this should result in the cation ratio of Sr:(Ti + V) > 1 for most of the prepared compositions. While strontium vacancies are known to be frequently encountered defects in SrTiO 3 -based solid solutions, the formation of vacancies in the titanium sublattice or the simultaneous formation of strontium and titanium vacancies were shown to be energetically unfavorable [33,[50][51][52]. Therefore, excess strontium can be compensated for by precipitation of a secondary Sr-rich phase such as Ruddlesden-Popper Sr 3 Ti 2 O 7 or strontium orthovanadate Sr 3 (VO 2 ) 2 . Given the stability of the latter phase and the difficulties faced during the synthesis, it would be reasonable to assume that traces of the Sr 3 (VO 2 ) 2 phase, undetected by XRD due to their very small amount and dispersed state, may still be present in the samples. In that case, the relationship between the nominal and actual compositions are described by the following equation: If Equation (1) is true, only one of the synthesized materials actually remains A-site deficient, while the fraction of precipitated Sr 3 (VO 4 ) 2 impurity should correspond to~2 mol.% in most cases (Table 2). Most of sintered STVN ceramics were dense, with a relative density between 93 and 97% of the theoretical density (roughly estimated from the structural data assuming the nominal cation composition) ( Table 1). The only exception was the S96V30N3 samples which showed a lower relative density of~89%. The estimated values of relative density are in good agreement with the microstructure of the STVN ceramics observed by SEM ( Figure S9). SEM images of fractured S96V30N3 ceramics revealed easily distinguishable individual grains with a size of 0.5-3.5 µm and visible porosity. Other compositions appeared to be comparatively dense, with a minor closed porosity and slightly different microstructural features. The S96V20N6 ceramics were composed of comparatively small grains (1.5-6.0 µm), but their shape implies an onset of formation of denser agglomerates. On the contrary, the SEM of fractured cross-sections of the S98V40N4 and S96V40N8 samples showed dense sintered bodies with mostly indistinguishable individual grains. The microstructure of the S100V20N2 and S98V20N4 samples is an intermediate case comprising both smaller easily distinguishable grains (2-10 µm) and large dense agglomerates with sizes up to 50 µm. Note that SEM/EDS analysis demonstrated a uniform distribution of elements (Sr, Ti, V, and O) in both small grains and larger dense blocks. The changes in microstructural features with composition are caused, apparently, by two main factors. First, increasing vanadium content appears to promote sinterability and grain growth. The second factor is the expected presence of the Sr 3 (VO 4 ) 2 secondary phase. Although the available phase diagram of the SrO-V 2 O 5 system predicts that Sr 3 (VO 4 ) 2 melts at 1545 • C [53,54], our previous experience of sintering of SrVO 3 -based ceramics [15,16] showed that an excess of Sr 3 (VO 4 ) 2 impurity promotes the melting of the samples already at 1500 • C in 10% H 2 -N 2 atmosphere even though the melting point of pure SrVO 3 should be as high as 1950-2000 • C [55,56]. Thus, the traces of Sr 3 (VO 4 ) 2 phase impurity ( Table 2) are likely to play the role of sintering aid, promoting the grain growth and densification, while the impurity-free S96V30N3 ceramics remain comparatively porous.
The distribution of metallic nickel in sintered STVN ceramics was inspected using SEM/EDS. Representative images are shown in Figure 3; the SEM/EDS micrographs for the entire range of compositions can be found in Figure S10. It was observed that the size of the Ni particles correlated with the grain size of the perovskite phase and, therefore, the density of grain boundaries (i.e., the number of available grain boundaries per unit volume). The samples with a denser microstructure and larger grains (e.g., S96V40N8) comprised larger Ni particles, with sizes up to~1 µm, while the ceramics with smaller well-defined individual grains of the perovskite phase had more evenly distributed smaller metallic Ni inclusions, with sizes varying from 0.2 to 0.7 µm (as in S100V20N2). Note that the easily visible agglomeration of the metallic nickel phase does not completely exclude the possibility of a partial incorporation of Ni cations into the B-sublattice of Sr(Ti,V)O 3-δ perovskites. microstructural features. The S96V20N6 ceramics were composed of comparatively small grains (1.5-6.0 μm), but their shape implies an onset of formation of denser agglomerates. On the contrary, the SEM of fractured cross-sections of the S98V40N4 and S96V40N8 samples showed dense sintered bodies with mostly indistinguishable individual grains. The microstructure of the S100V20N2 and S98V20N4 samples is an intermediate case comprising both smaller easily distinguishable grains (2-10 μm) and large dense agglomerates with sizes up to 50 μm. Note that SEM/EDS analysis demonstrated a uniform distribution of elements (Sr, Ti, V, and O) in both small grains and larger dense blocks. The changes in microstructural features with composition are caused, apparently, by two main factors. First, increasing vanadium content appears to promote sinterability and grain growth. The second factor is the expected presence of the Sr3(VO4)2 secondary phase. Although the available phase diagram of the SrO-V2O5 system predicts that Sr3(VO4)2 melts at 1545 °C [53,54], our previous experience of sintering of SrVO3-based ceramics [15,16] showed that an excess of Sr3(VO4)2 impurity promotes the melting of the samples already at 1500 °C in 10% H2-N2 atmosphere even though the melting point of pure SrVO3 should be as high as 1950-2000 °C [55,56]. Thus, the traces of Sr3(VO4)2 phase impurity ( Table 2) are likely to play the role of sintering aid, promoting the grain growth and densification, while the impurity-free S96V30N3 ceramics remain comparatively porous.
The distribution of metallic nickel in sintered STVN ceramics was inspected using SEM/EDS. Representative images are shown in Figure 3; the SEM/EDS micrographs for the entire range of compositions can be found in Figure S10. It was observed that the size of the Ni particles correlated with the grain size of the perovskite phase and, therefore, the density of grain boundaries (i.e., the number of available grain boundaries per unit volume). The samples with a denser microstructure and larger grains (e.g., S96V40N8) comprised larger Ni particles, with sizes up to ~1 μm, while the ceramics with smaller well-defined individual grains of the perovskite phase had more evenly distributed smaller metallic Ni inclusions, with sizes varying from 0.2 to 0.7 μm (as in S100V20N2). Note that the easily visible agglomeration of the metallic nickel phase does not completely exclude the possibility of a partial incorporation of Ni cations into the B-sublattice of Sr(Ti,V)O3-δ perovskites.

Thermomechanical Compatibility of STVN with Solid Electrolyte
The dilatometric curves of the STVN ceramics showed smooth, slightly non-linear behavior upon heating in a reducing 10% H2-N2 atmosphere. The representative examples for the selected samples are depicted in Figure 4. Similar to SrTi1-yVyO3-δ series [34] and other SrVO3-δ-based materials [14,15,22], increasing deviation from the linear thermal

Thermomechanical Compatibility of STVN with Solid Electrolyte
The dilatometric curves of the STVN ceramics showed smooth, slightly non-linear behavior upon heating in a reducing 10% H 2 -N 2 atmosphere. The representative examples for the selected samples are depicted in Figure 4. Similar to SrTi 1-y V y O 3-δ series [34] and other SrVO 3-δ -based materials [14,15,22], increasing deviation from the linear thermal expansion at higher temperatures can be partly attributed to a minor contribution of chemical expansion associated with the variable oxygen content in the perovskite lattice.  The average linear thermal expansion coefficients (TECs) of the STVN ceramics calculated from the dilatometric data were found to vary in the range of 12.5-14.3 ppm/K at 25-1100 °C, generally decreasing with an increase in titanium content in the perovskite lattice (Table 3). This is consistent with the observation reported earlier for SrVO3-SrTiO3 solid solutions [34] where increasing titanium content was found to suppress both true thermal expansion and chemical strain and to reduce the average TEC (100-1000 °C) from 19.3 ppm/K for undoped SrVO3-δ to 12.0 ppm/K for titanium-rich SrTi0.9V0.1O3-δ ceramics. 10.0 ± 0.1 3.58 ± 0.02 1 σ = A0 × exp(−EA/(RT)) where EA is the activation energy and A0 is the pre-exponential factor.
The TECs of STVN ceramics slightly exceed the corresponding value for the common 8YSZ solid electrolyte ceramics (10.5 ppm/K at 25-1100 °C [57]) but are still very close. This should ensure good thermomechanical compatibility between porous STVN electrodes and 8YSZ solid electrolytes during thermal cycling.

Electrical Conductivity
Sintered STVN ceramics exhibit semiconducting behavior with temperatureactivated electrical conductivity under reducing conditions ( Figure 5A). Electrical conductivity was found to increase with increasing vanadium content in the perovskite lattice, reaching a maximum of ~17 S/cm at 900 °C for the S98V40Ni4 composition, while the activation energy for conductivity showed an opposite trend (Table 3). These observations are in good agreement with previously reported data on the SrVO3-SrTiO3 pseudobinary system [34,48]. SrVO3-δ is known to be a metallic conductor with a The average linear thermal expansion coefficients (TECs) of the STVN ceramics calculated from the dilatometric data were found to vary in the range of 12.5-14.3 ppm/K at 25-1100 • C, generally decreasing with an increase in titanium content in the perovskite lattice (Table 3). This is consistent with the observation reported earlier for SrVO 3 -SrTiO 3 solid solutions [34] where increasing titanium content was found to suppress both true thermal expansion and chemical strain and to reduce the average TEC (100-1000 • C) from 19.3 ppm/K for undoped SrVO 3-δ to 12.0 ppm/K for titanium-rich SrTi 0.9 V 0.1 O 3-δ ceramics.
where E A is the activation energy and A 0 is the pre-exponential factor.
The TECs of STVN ceramics slightly exceed the corresponding value for the common 8YSZ solid electrolyte ceramics (10.5 ppm/K at 25-1100 • C [57]) but are still very close. This should ensure good thermomechanical compatibility between porous STVN electrodes and 8YSZ solid electrolytes during thermal cycling.

Electrical Conductivity
Sintered STVN ceramics exhibit semiconducting behavior with temperature-activated electrical conductivity under reducing conditions ( Figure 5A). Electrical conductivity was found to increase with increasing vanadium content in the perovskite lattice, reaching a maximum of~17 S/cm at 900 • C for the S98V40Ni4 composition, while the activation energy for conductivity showed an opposite trend (Table 3). These observations are in good agreement with previously reported data on the SrVO 3 -SrTiO 3 pseudobinary system [34,48]. SrVO 3-δ is known to be a metallic conductor with a conductivity level as high as 10 3 S/cm at 700-900 • C [12,16]. On the other hand, undoped SrTiO 3-δ is a wide-bandgap semiconductor with a conductivity that is ≤10 −1 S/cm at 900 • C under similar reducing conditions [23][24][25]. Accordingly, decreasing the vanadium content in SrTi 1-y V y O 3-δ solid solutions was demonstrated to result in a gradual decrease in electrical conductivity and a transition from metallic to semiconducting behavior [34,48]. This was interpreted in terms of a gradual change in the mechanism of electronic conduction. conductivity level as high as ~10 3 S/cm at 700-900 °C [12,16]. On the other hand, undoped SrTiO3-δ is a wide-bandgap semiconductor with a conductivity that is ≤10 −1 S/cm at 900 °C under similar reducing conditions [23][24][25]. Accordingly, decreasing the vanadium content in SrTi1-yVyO3-δ solid solutions was demonstrated to result in a gradual decrease in electrical conductivity and a transition from metallic to semiconducting behavior [34,48]. This was interpreted in terms of a gradual change in the mechanism of electronic conduction. The charged point defects in Sr1-xTi1-yVyO3-δ solid solutions based on the Sr 2+ Ti 4+ O 2− host lattice include vacancies in the strontium sublattice, Ti 3+ and V 3+ cations in the Bsublattice, and vacancies in the oxygen sublattice. The corresponding electroneutrality condition is given by (using Kröger-Vink notation): where ▯ indicates a vacancy, Ti ≡ Ti 3+ , and V ≡ V 3+ . It is expected that the electronic transport in the compositions with a vanadium content of ~20-40 at.% in the titanium sublattice is carried out by the electrons localized on the titanium cations and is contributed to by electron hopping between V 4+ /V 3+ pairs [34]. Thus, the concentration of electronic charge carriers is equivalent to the concentration of Ti 3+ and Vi 3+ . As vanadium cations are more easily reducible, the electronic transport via V 4+ /V 3+ redox pairs should prevail and declines when the total vanadium content decreases. The variations in the conductivity of the samples with a similar vanadium content in the perovskite lattice can be attributed to an interplay between the porosity, microstructural features, and the distribution of insulating Sr3(VO4)2 impurity. Reducing oxygen partial pressure results in a modest increase in the conductivity for all STVN compositions ( Figure 5B). This should be attributed to an increase in the concentration of n-type electronic charge carriers due to the reduction of B-site transition metal cations and oxygen release from the lattice: where B is Ti or V. The dependence of conductivity on p(O2) is very weak and resembles that reported for titanium-rich SrTi1-yVyO3-δ solid solutions under similar conditions [34]. The electrical conductivity of electrode materials for SOFC/SOEC should be as high as possible to minimize ohmic losses and avoid limiting effects on electrochemical activity. It is generally considered that the minimum required electronic conductivity corresponds The charged point defects in Sr 1-x Ti 1-y V y O 3-δ solid solutions based on the Sr 2+ Ti 4+ O 2− host lattice include vacancies in the strontium sublattice, Ti 3+ and V 3+ cations in the Bsublattice, and vacancies in the oxygen sublattice. The corresponding electroneutrality condition is given by (using Kröger-Vink notation): where indicates a vacancy, Ti Ti ≡ Ti 3+ , and V Ti ≡ V 3+ . It is expected that the electronic transport in the compositions with a vanadium content of~20-40 at.% in the titanium sublattice is carried out by the electrons localized on the titanium cations and is contributed to by electron hopping between V 4+ /V 3+ pairs [34]. Thus, the concentration of electronic charge carriers is equivalent to the concentration of Ti 3+ and Vi 3+ . As vanadium cations are more easily reducible, the electronic transport via V 4+ /V 3+ redox pairs should prevail and declines when the total vanadium content decreases. The variations in the conductivity of the samples with a similar vanadium content in the perovskite lattice can be attributed to an interplay between the porosity, microstructural features, and the distribution of insulating Sr 3 (VO 4 ) 2 impurity. Reducing oxygen partial pressure results in a modest increase in the conductivity for all STVN compositions ( Figure 5B). This should be attributed to an increase in the concentration of n-type electronic charge carriers due to the reduction of B-site transition metal cations and oxygen release from the lattice: where B is Ti or V. The dependence of conductivity on p(O 2 ) is very weak and resembles that reported for titanium-rich SrTi 1-y V y O 3-δ solid solutions under similar conditions [34]. The electrical conductivity of electrode materials for SOFC/SOEC should be as high as possible to minimize ohmic losses and avoid limiting effects on electrochemical activity. It is generally considered that the minimum required electronic conductivity corresponds to >1 S/cm for porous electrode structures and one order of magnitude higher value for the bulk materials [58]. The conductivity of the STVN ceramics with the highest vanadium content, S98V40N4 and S96V40N8, amounts to 10-17 S/cm at 700-900 • C and, thus, appears to be acceptable for electrode application. Other STVN materials exhibit 5-12 times lower electrical conductivity in the same temperature range. Hence, comparatively low conductivity may be one of the performance-limiting factors in the case of electrode structures based on these compositions.

As-Prepared STVN Electrodes
The initial electrochemical experiments aimed at a comparative assessment of the activity of plain as-prepared STVN porous electrode layers applied onto 8YSZ solid electrolytes using a symmetrical electrode|electrolyte|electrode cell configuration. to >1 S/cm for porous electrode structures and one order of magnitude higher value for the bulk materials [58]. The conductivity of the STVN ceramics with the highest vanadium content, S98V40N4 and S96V40N8, amounts to 10-17 S/cm at 700-900 °C and, thus, appears to be acceptable for electrode application. Other STVN materials exhibit 5-12 times lower electrical conductivity in the same temperature range. Hence, comparatively low conductivity may be one of the performance-limiting factors in the case of electrode structures based on these compositions.

As-Prepared STVN Electrodes
The initial electrochemical experiments aimed at a comparative assessment of the activity of plain as-prepared STVN porous electrode layers applied onto 8YSZ solid electrolytes using a symmetrical electrode│electrolyte│electrode cell configuration.  A typical impedance spectrum of a symmetrical cell with STVN electrodes is presented in Figure 7A (Nyquist plot) and Figure 7B (Bode plot). All collected spectra were satisfactorily fitted to a simple equivalent circuit (inset in Figure 7A) comprising an inductive element (L), the ohmic resistance of the cell (R Ohm ), and two (R||CPE) contributions (where R is resistance and CPE is a constant phase element [59,60]) at high and low frequencies (HF and LF, respectively) corresponding to the electrode processes. The sum of R HF and R LF provides the specific electrode polarization resistance (R p ). Taking into account the literature data (e.g., Refs. [61,62]), the HF contribution (frequencies in the order of kHz) is presumably associated with the ionic/electronic charge transfer at the triple-phase boundary (TPB), whereas LF components (frequencies below 1 kHz) are generally ascribed to the gas kinetics at the electrode surface (non-Faradaic mass transport mechanisms such as hydrogen adsorption/desorption and surface diffusion of active species). A typical impedance spectrum of a symmetrical cell with STVN electrodes is presented in Figure 7A (Nyquist plot) and Figure 7B (Bode plot). All collected spectra were satisfactorily fitted to a simple equivalent circuit (inset in Figure 7A) comprising an inductive element (L), the ohmic resistance of the cell (ROhm), and two (R||CPE) contributions (where R is resistance and CPE is a constant phase element [59,60]) at high and low frequencies (HF and LF, respectively) corresponding to the electrode processes. The sum of RHF and RLF provides the specific electrode polarization resistance (Rp). Taking into account the literature data (e.g., Refs. [61,62]), the HF contribution (frequencies in the order of kHz) is presumably associated with the ionic/electronic charge transfer at the triple-phase boundary (TPB), whereas LF components (frequencies below 1 kHz) are generally ascribed to the gas kinetics at the electrode surface (non-Faradaic mass transport mechanisms such as hydrogen adsorption/desorption and surface diffusion of active species). The STVN electrodes were found to exhibit rather poor electrocatalytic activity for hydrogen oxidation reaction even at 900 °C, with Rp values varying between 9.8 and 26.4 Ohm × cm 2 ( Figures 8A). In all cases, the LF contribution was the dominant performancelimiting process. Decreasing temperature results in slower kinetics of the electrode process, with both HF and LF contributions exponentially increasing upon cooling. Figure  8B illustrates the evolution of EIS spectra between 700 and 900 °C for the S96V30N3-based cells, with the Rp value increasing from 9.8 to 94 Ohm × cm 2 upon cooling. The temperature dependence of electrode polarization resistance is depicted in Figure 9. The S96V30N3 and S96V20N6 electrodes exhibit the best and the worst performance, respectively, in the entire temperature range. At the same time, the values of EA vary in a wide range from 95 kJ/mol (S100V20N2) to 141 kJ/mol (S98V40N4). The STVN electrodes were found to exhibit rather poor electrocatalytic activity for hydrogen oxidation reaction even at 900 • C, with R p values varying between 9.8 and 26.4 Ohm × cm 2 ( Figure 8A). In all cases, the LF contribution was the dominant performance-limiting process. Decreasing temperature results in slower kinetics of the electrode process, with both HF and LF contributions exponentially increasing upon cooling. Figure 8B illustrates the evolution of EIS spectra between 700 and 900 • C for the S96V30N3-based cells, with the R p value increasing from 9.8 to 94 Ohm × cm 2 upon cooling. The temperature dependence of electrode polarization resistance is depicted in Figure 9. The S96V30N3 and S96V20N6 electrodes exhibit the best and the worst performance, respectively, in the entire temperature range. At the same time, the values of E A vary in a wide range from 95 kJ/mol (S100V20N2) to 141 kJ/mol (S98V40N4).
attributed to the poor intrinsic electrocatalytic activity of Sr1-xTi1-yVyO3-δ perovskites in combination with a low expected ionic conductivity in these phases. The high firing temperature required to eliminate an undesired insulating Sr3(VO4)2 impurity phase promoted the agglomeration of catalytically active Ni particles. Therefore, it can be assumed that the differences in the obtained Rp values for different STVN electrodes are mainly caused by a somewhat different uncontrolled distribution of the segregated nickel phase.   unsatisfactory electrochemical performance of the as-prepared STVN electrodes can be attributed to the poor intrinsic electrocatalytic activity of Sr1-xTi1-yVyO3-δ perovskites in combination with a low expected ionic conductivity in these phases. The high firing temperature required to eliminate an undesired insulating Sr3(VO4)2 impurity phase promoted the agglomeration of catalytically active Ni particles. Therefore, it can be assumed that the differences in the obtained Rp values for different STVN electrodes are mainly caused by a somewhat different uncontrolled distribution of the segregated nickel phase.   The analysis of the results did not reveal any evident correlation between the nominal nickel content in the STVN, the electrical conductivity of ceramics, electrode polarization resistance, and the corresponding activation energy. Therefore, one may conclude that the unsatisfactory electrochemical performance of the as-prepared STVN electrodes can be attributed to the poor intrinsic electrocatalytic activity of Sr 1-x Ti 1-y V y O 3-δ perovskites in combination with a low expected ionic conductivity in these phases. The high firing temperature required to eliminate an undesired insulating Sr 3 (VO 4 ) 2 impurity phase promoted the agglomeration of catalytically active Ni particles. Therefore, it can be assumed that the differences in the obtained R p values for different STVN electrodes are mainly caused by a somewhat different uncontrolled distribution of the segregated nickel phase.

Chemical Reactivity between STVN and Solid Electrolytes
The reactivity between the electrode and the electrolyte material during the fabrication and/or operation of solid oxide cells is one of the factors that may have a negative impact on the overall performance of SOFCs. The undesired cation interdiffusion may result in the formation of insulating products at the electrode/electrolyte interface, blocking the mass-transfer processes (oxygen diffusion). In the present study, the SEM/EDS inspection of the electrode/electrolyte assemblies after fabrication and after experiments did not reveal a formation of any interlayers at the STVN/8YSZ interfaces. Since the thickness of such layers may be too low for the resolution of the SEM/EDS, additional reactivity tests were conducted. Powdered STVN + 8YSZ mixtures (50:50 wt.%) were fired at 1200 • C for 10 h in a flowing 10% H 2 -N 2 atmosphere and then analyzed using XRD. The appearance of a very small additional peak at 2Θ ≈ 30.9 • was detected in the XRD patterns after calcination ( Figure 10). This peak can be assigned to the perovskite-like SrZrO 3 phase, a poorly conducting reactivity product that is frequently observed at the (La,Sr)MnO 3 /8YSZ interface and has a negative effect on the performance of (La,Sr)MnO 3 -based cathodes in SOFCs [63]. Nonetheless, the reactivity between 8YSZ and STVN under applied firing conditions was found to be very limited. As the actual STVN electrode layers in the present work were sintered at a lower temperature (1100 • C) and for a shorter time (2 h), one may assume a rather limited adverse impact of possible reactivity between the materials on the electrode performance.

Modification of STVN Electrodes via Infiltration
As plain STVN electrodes demonstrated poor electrochemical activity, the modification of porous electrode layers through the incorporation of additional components-CGO as the oxygen-ion-conducting phase and nano-sized Ni as the A common practice to improve the electrochemical activity of electrode materials with predominant electronic conductivity and poor ionic transport is to introduce a fraction of solid electrolyte material into the porous electrode layer, thus forming a composite electrode. This results in a drastic increase in the triple-phase boundary length (i.e., the concentration of sites where the electrochemical reaction takes place) and, correspondingly, an improvement in electrochemical activity. As gadolinia-doped ceria (CGO) is often used for this purpose, the reactivity between CGO (Ce 0.9 Gd 0.1 O 2-δ in the present work) and STVN was tested in a similar manner as in the case for 8YSZ. A careful inspection of the XRD patterns of the STVN + CGO mixtures (50:50 wt.%) after calcination at 1200 • C for 10 h in 10% H 2 -N 2 atmosphere revealed the presence of several additional very small peaks on the background level ( Figure 10). These tiny peaks coincided with the positions of the main reflections of the Sr 3 (VO 4 ) 2 phase. As discussed above, the traces of this phase are expected to be present in the synthesized STVN to compensate for the segregation of metallic nickel. It is possible that a partial reduction of CGO at 1200 • C accompanied by the release of oxygen from the fluorite lattice may promote the transformation of an additional fraction of Sr(Ti,V)O 3-δ perovskite into the oxidized strontium orthovanadate phase. However, no evidence of chemical reactivity between STVN and CGO with the formation of additional phases could be detected at temperatures as high as 1200 • C.

Modification of STVN Electrodes via Infiltration
As plain STVN electrodes demonstrated poor electrochemical activity, the modification of porous electrode layers through the incorporation of additional components-CGO as the oxygen-ion-conducting phase and nano-sized Ni as the electrocatalyst-was attempted in order to improve their performance. Two STVN compositions were selected for these studies: one with the best performance in the unmodified form (S96V30N3) and one with the average performance (S98V40N4). The modification was completed via the infiltration procedure described in Section 2.4. Three combinations of additional components were tested including (i) only CGO, (ii) CGO + Ni with a molar ratio 10:1, and (iii) CGO + Ni with a molar ratio 2:1. The total load of infiltrated components after thermal treatment corresponded to 21-30 wt.%. This provides the fraction of extra Ni catalyst equal tõ 1 and~4 wt.% in the case of the CGO:Ni ratio of 10:1 and 2:1, respectively. The SEM/EDS inspection of modified electrodes revealed that the porous STVN backbone was uniformly covered by submicron CGO particles (100-150 nm) and well-dispersed nanostructured Ni catalysts ( Figure 11, Figures S11 and S12). The incorporation of CGO (~21 wt.%) as a solid electrolyte component resulted in a reduction in electrode polarization resistance ( Figure 12). While this change was moderate in the case of the S96V30N3-based cell, from 9.8 to 6.6 Ohm × cm 2 at 900 °C, the Rp of the S98V40N4 electrodes dropped more than four times from ~18 to 4.2 Ohm × cm 2 at 900 °C. The incorporation of CGO (~21 wt.%) as a solid electrolyte component resulted in a reduction in electrode polarization resistance ( Figure 12). While this change was moderate in the case of the S96V30N3-based cell, from 9.8 to 6.6 Ohm × cm 2 at 900 • C, the R p of the S98V40N4 electrodes dropped more than four times from~18 to 4.2 Ohm × cm 2 at 900 • C. The enhanced performance is attributed to an increase in the TPB length as a result of the introduction of the oxygen-ion conducting phase into the electrode structure.
Co-infiltration of CGO and Ni (27-30 wt.%, CGO:Ni = 10:1) substantially improved the electrochemical activity of STVN electrodes; the R p values dropped 5-7 times compared with the as-prepared electrodes. Again, this enhancement was stronger for the S98V40N4-based cell, where electrode polarization resistance decreased to 2.3 Ohm × cm 2 at 900 • C. Compared with the electrodes infiltrated only with CGO, the decrease in R p values for electrodes co-modified with CGO + Ni is mainly associated with the reduced LF contribution. This implies that the addition of a dispersed Ni catalyst leads to enhanced kinetics of processes at the electrode/gas interface. The incorporation of CGO (~21 wt.%) as a solid electrolyte component resulted in a reduction in electrode polarization resistance ( Figure 12). While this change was moderate in the case of the S96V30N3-based cell, from 9.8 to 6.6 Ohm × cm 2 at 900 °C, the Rp of the S98V40N4 electrodes dropped more than four times from ~18 to 4.2 Ohm × cm 2 at 900 °C. The enhanced performance is attributed to an increase in the TPB length as a result of the introduction of the oxygen-ion conducting phase into the electrode structure. Co-infiltration of CGO and Ni (27-30 wt.%, CGO:Ni = 10:1) substantially improved the electrochemical activity of STVN electrodes; the Rp values dropped 5-7 times compared with the as-prepared electrodes. Again, this enhancement was stronger for the S98V40N4-based cell, where electrode polarization resistance decreased to 2.3 Ohm × cm 2 at 900 °C. Compared with the electrodes infiltrated only with CGO, the decrease in Rp values for electrodes co-modified with CGO + Ni is mainly associated with the reduced Figure 12. Impedance spectra of as-prepared (unmodified) and infiltrated S98V40N4 (A) and S96V30N3 (B) electrodes in humidified 10% H 2 -N 2 at 900 • C. All spectra are corrected for ohmic contributions. Insets show magnified sections of the plots.
Further increasing the fraction of infiltrated nickel (CGO:Ni = 2:1) was found to have a comparatively small effect on the electrode polarization resistance of S96V30N3-based electrodes (R p = 1.6 Ohm × cm 2 at 900 • C) or even a slightly negative impact in the case of S98V40N4 (2.6 Ohm × cm 2 at 900 • C). This seems to indicate that an increased Ni load promotes the agglomeration of Ni particles rather than their better distribution. As a result, the concentration of catalytically active sites and the electrode kinetics are not further improved. Another observation is the emergence of an additional contribution at low frequencies (f < 3 Hz) in the impedance spectra of the infiltrated cells. According to the literature, this contribution can be assigned to the gas diffusion processes. As these processes are mainly affected by the electrode microstructure, the magnitude of this contribution may vary from one electrode to another and does not depend on temperature [64]. In the present case, it becomes visible in the spectra when the overall polarization resistance decreases and the infiltrated components reduce the porosity of the electrode layer, apparently introducing gas-diffusion limitations. Figure 13 presents the temperature dependence of electrode polarization resistance of unmodified and infiltrated electrodes. The incorporation of CGO or CGO in combination with Ni gradually decreases the activation energy of the electrode process (Table 4), thus leading to a stronger relative improvement in overall electrode performance at lower temperatures. Interestingly, both E A and R p values for S98V40N4 and S96V30N3 coinfiltrated with CGO and Ni (ratio 10:1) become similar. This seems to confirm that the porous STVN skeleton plays a rather passive role in terms of the electrocatalytic activity of infiltrated electrode layers.
layer, apparently introducing gas-diffusion limitations. Figure 13 presents the temperature dependence of electrode polarizat of unmodified and infiltrated electrodes. The incorporation of CGO combination with Ni gradually decreases the activation energy of the elec (Table 4), thus leading to a stronger relative improvement in ove performance at lower temperatures. Interestingly, both EA and Rp values and S96V30N3 co-infiltrated with CGO and Ni (ratio 10:1) become similar. confirm that the porous STVN skeleton plays a rather passive role in electrocatalytic activity of infiltrated electrode layers. Figure 13. Temperature dependence of the area-specific electrode polarization r unmodified and infiltrated S98V40N4 and S96V30N3 electrodes in humidified 10% Figure 13. Temperature dependence of the area-specific electrode polarization resistance R p of unmodified and infiltrated S98V40N4 and S96V30N3 electrodes in humidified 10% H 2 -N 2 .

Conclusions
A series of strontium titanate-vanadate (STVN) solid solutions with the nominal cation composition of Sr 1-x Ti 1-y-z V y Ni z O 3-δ (x = 0-0.04, y = 0.20-0.40 and z = 0.02-0.12) were prepared by the solid-state reaction route in reducing 10% H 2 -N 2 atmosphere. High-energy mechanochemical pre-treatments of the precursor mixtures have a weak effect in terms of the formation of the target perovskite phase. Prolonged firing at 1200 • C is required to eliminate the highly stable Sr 3 (VO 4 ) 2 intermediate phase. Under such conditions, Ni tends to segregate as a metallic phase and is unlikely to incorporate into the perovskite structure. The electrical conductivity of ceramics sintered at 1500 • C is temperature-activated and increases with increasing vanadium content reaching a maximum of~17 S/cm at 1000 • C. The prepared STVN materials show good thermomechanical and chemical compatibility with 8YSZ and CGO solid electrolytes. Porous STVN electrodes applied onto 8YSZ solid electrolytes demonstrate rather poor electrochemical activity in diluted H 2 . This is attributed to the low intrinsic electrocatalytic activity and low ionic conductivity of Sr(Ti,V)O 3-δ perovskites as well as the agglomeration of metallic Ni during the synthetic procedure. The overall performance of STVN electrodes can be substantially improved by the additions of CGO as an oxygen-ion conducting phase (20-30 wt.%) and surface modification by nano-sized nickel as an electrocatalyst (≥1 wt.%).
The preparation of single-phase Sr 1-x Ti 1-y-z V y Ni z O 3-δ perovskites and fabrication of Sr(Ti,V)O 3 -based electrodes decorated with a nanostructured Ni catalyst via an exsolution process requires the development of an alternative synthetic procedure which will make it possible to eliminate or avoid the formation of undesired phases. This is planned as a continuation of the present work. Funding: This study was completed in the scope of the projects CARBOSTEAM (POCI-01-0145-FEDER-032295) and HEALING (POCI-01-0145-FEDER-032036) funded by FEDER through COM-PETE2020-Programa Operacional Competitividade e Internacionalização (POCI) and national funds through FCT/MCTES and the project CICECO-Aveiro Institute of Materials (UIDB/50011/2020 and UIDP/50011/2020) financed by national funds through the FCT/MCTES and when appropriate co-financed by FEDER under the PT2020 Partnership Agreement.

Data Availability Statement:
Data is contained within the article and/or available from the corresponding author upon reasonable request.